Forged aluminum alloy having excellent strength and ductility and method for producing the same

ABSTRACT

Provided is a hot-forged 6xxx-series aluminum alloy having excellent corrosion resistance and still having both high strength and good ductility. A forged 6xxx-series aluminum alloy having a specific chemical composition after solution treatment is further subjected to warm working to introduce dislocations into the forged aluminum alloy microstructure. This allows the forged aluminum alloy after artificial aging to have a microstructure which has a high dislocation density, includes a large proportion of small angle grain boundaries, and has a high average number density of precipitates. Thus, the resulting forged aluminum alloy has a 0.2% yield strength of 400 MPa or more and an elongation of 10% or more and combines properties necessary for suspension parts.

FIELD OF INVENTION

The present invention relates to a forged aluminum alloy having excellent strength and ductility, and a method for producing the same. Hereinafter, “aluminum” is also simply referred to as “AI”.

As used herein, the term “forged material” refers to a forged aluminum alloy produced (plastically worked) by hot forging.

In the present the present invention, the term “forged material” is used not as a term describing a production process of a product, but as a term which is well-known to be generally used as a technical term and/or a patent term for specifying the state of the product.

Aluminum alloy materials, when having different plastic working histories as in hot-forged materials, extruded materials (extrusions), and rolled materials, quite differ from each other in microstructures and properties, even when having identical alloy chemical compositions. Thus, specifying or defining of the alloy chemical compositions, microstructure, and properties of an aluminum alloy material has no point unless the plastic working history of the aluminum alloy material is specified.

Accordingly, the term “hot-forged material” or a synonym thereof “forged material” is used in the appended claims and in the following description so as to clearly distinguish a target hot-forged aluminum alloy, to which the present invention is applied, from other plastically worked materials and to clearly specify the state of the substance (material).

BACKGROUND OF INVENTION

Reduction in body weight of and resulting improvements in fuel efficiency of, automobiles and other transportation equipment have been pursued so as to counter global environmental issues caused typically by exhaust gases. To this end, 6xxx-series (Al—Mg—Si) hot-forged aluminum alloys as prescribed in Aluminum Association (AA) standards or Japanese Industrial Standards (JIS) are used for structural components and structural parts of automobiles and other transportation equipment and, in particular, for automobile suspension parts such as upper arms and lower arms.

Hot-forged 6xxx-series aluminum alloys, when used for these structural components and structural parts, offer high strength and high toughness and have relatively excellent corrosion resistance. Hereinafter, such structural components and structural parts of transportation equipment will be illustrated by taking automobile suspension parts as an example.

For further weight reduction of automobiles, automobile suspension parts require thinner thicknesses and still require higher strength and higher toughness. The automobile suspension parts also function as safety-related parts and require higher corrosion resistance to intergranular corrosion (grain-boundary corrosion) and to stress corrosion cracking so as to ensure reliability as the safety-related parts. Accordingly, various techniques have been developed to improve chemical compositions and microstructures of material hot-forged 6xxx-series aluminum alloys.

For example, Japanese Patent No. 5110938 proposes a technique, in which a forged 6xxx-series aluminum alloy is controlled to have a precipitate density of 1.5% or less in terms of average area percentage, and grain-boundary precipitates are controlled to be present at an average spacing between the precipitates of 0.7 μm or more, where the grain-boundary precipitates are observed in a microstructure in a cross-sectional region including a parting line formed upon forging.

Japanese Patent No. 5723192 proposes another technique. This technique relates to a forged 6xxx-series aluminum alloy formed by subjecting an aluminum alloy extrusion to hot forging. The aluminum alloy has an unrecrystallized region in the entire cross section thereof. The unrecrystallized region includes small angle grain boundaries with a tilt angle of from 2° to less than 15°, and large angle grain boundaries with a tilt angle of 15° or more. In the unrecrystallized region, areas surrounded by boundaries with a tilt angle of 2° or more have an average grain size of 10 μm or less. The unrecrystallized region occupies 75% or more of the entire cross section. The unrecrystallized structure region includes dispersed particles having a maximum length of 10 nm to 800 nm in an average number density of 10 per cubic micrometer. The unrecrystallized structure region includes precipitates having a maximum length of 0.5 μm or more in an average area percentage of 2.5% or less.

In contrast, though not in the field of hot-forged aluminum alloys, Japanese Unexamined Patent Application Publication (JP-A) No. 2014-218685 and Japanese Patent No. 5082483 each disclose a metallurgical technique so as to offer higher strength of aluminum alloy materials. With this known technique, a 6xxx-series aluminum alloy ingot is sequentially subjected to solution treatment (solution heat treatment), repeatedly to warm forging at about 150° C. to about 250° C., and to artificial aging.

SUMMARY OF INVENTION

However, even the improvements in chemical compositions and microstructures of hot-forged 6xxx-series aluminum alloys, as disclosed typically in Japanese Patent No. 5110938 and Japanese Patent No. 5723192, are susceptible to improvement so as to give forged aluminum alloys having strength and ductility both at excellent levels, where the strength and ductility properties are mutually contradictory and resist being compatible with each other.

With the technique of subjecting a 6xxx-series aluminum alloy ingot repeatedly to warm forging and then to artificial aging so as to offer higher strength as disclosed typically in JP-A No. 2014-218685 and Japanese Patent No. 5082483, it has been considered that the aluminum alloy ingot, when subjected to hot forging at a high temperature of typically 500° C. less effectively has such higher strength. It is still unknown that this technique is also effective for better mechanical properties of hot-forged 6xxx-series aluminum alloys.

The present invention has been made while focusing on these circumstances and has an object to provide a forged 6xxx-series aluminum alloy that has, as a precondition, excellent corrosion resistance and still has strength and ductility both at excellent levels (has both high strength and high elongation).

To achieve the object, the present invention provides a forged aluminum alloy having excellent strength and ductility. The forged aluminum alloy contains, in mass percent, Si in a content of 0.7% to 1.5%, Mg in a content of 0.6% to 1.2%, Fe in a content of 0.01% to 0.5%, and at least one element selected from the group consisting of Mn in a content of 0.05% to 1.0%, Cr in a content of 0.01% to 0.5%, and Zr in a content of 0.01% to 0.2%, with the remainder consisting of Al and inevitable impurities. The forged aluminum alloy has a microstructure in an observation plane at a center of the thickness in a thickest portion of the forged aluminum alloy. The microstructure has a dislocation density of from 1.0×10¹⁴ to 5.0×10¹⁶ per square meter on average as measured by X-ray diffractometry. The microstructure includes small angle grain boundaries with a tilt angle of 2° to 15° in an average proportion of 50% or more as measured by SEM-EBSD analysis, where the small angle grain boundaries are present around grains having a misorientation of 2° or more. The microstructure includes precipitates measurable with a transmission electron microscope (TEM) at 300000-fold magnification in an average number density of 5.0×10² per cubic micrometer or more.

To achieve the object, the present invention also provides a method for producing a forged aluminum alloy having excellent strength and ductility. The method includes preparing an aluminum alloy ingot. The aluminum alloy ingot contains, in mass percent, Si in a content of 0.7% to 1.5%, Mg in a content of 0.6% to 1.2%, Fe in a content of 0.01% to 0.5%, and at least one element selected from the group consisting of Mn in a content of 0.05% to 1.0%, Cr in a content of 0.01% to 0.5%, and Zr in a content of 0.01% to 0.2%, with the remainder consisting of Al and inevitable impurities. The aluminum alloy ingot is subjected sequentially to homogenization and hot forging to give a forged material. The forged material is further subjected sequentially to solution treatment, quenching, warm working, and artificial aging in the specified sequence. The forged material (forged aluminum alloy) after artificial aging has a microstructure in an observation plane at a center of the thickness in a thickest portion of the forged aluminum alloy. The microstructure has a dislocation density of from 1.0×10¹⁴ to 5.0×10¹⁶ per square meter on average as measured by X-ray diffractometry. The microstructure includes small angle grain boundaries with a tilt angle of 20 to 15° in an average proportion of 50% or more as measured by SEM-EBSD analysis, where the small angle grain boundaries are present around grains having a misorientation of 20 or more. The microstructure includes precipitates measurable with a transmission electron microscope (TEM) at 300000-fold magnification in an average number density of 5.0×10² per cubic micrometer or more.

It has been found in the present invention that, when a forged 6xxx-series aluminum alloy after solution treatment and quenching is subjected to warm working to be imparted with working strain, and is then subjected to artificial aging, the resulting forged aluminum alloy has strength and ductility both at higher levels (has both higher strength and better ductility) as compared with regular equivalents to which working stain is not imparted.

To offer or to ensure the advantageous effect, the present invention specifies a microstructure in a central part of the thickness in a thickest portion of the forged material after artificial aging. Specifically, the present invention specifies the average dislocation density, the average proportion of small angle grain boundaries, and the average number density of precipitates in the microstructure, as described above.

The present invention allows forged 6xxx-series aluminum alloys to have strength and ductility both at higher levels, and this enables further reduction in weight.

DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

Some embodiments of the present invention will be illustrated below.

Chemical Composition

Initially, the chemical composition of an aluminum alloy constituting an ingot as a material for the forged material (forged aluminum alloy) according to the present invention and for an aluminum alloy ingot will be described below.

The chemical composition of the 6xxx-series (Al—Mg—Si-series) aluminum alloy for use in the present invention should be determined or specified so as to ensure higher strength better ductility, and high corrosion resistance or durability to be used typically as the forged suspension parts. Within the range of 6xxx-series aluminum alloy chemical compositions, the aluminum alloy for use in the present invention contains, in a chemical composition in mass percent Si in a content of 0.7% to 1.5%, Mg in a content of 0.6% to 1.2%, Fe in a content of 0.01% to 0.5%, and at least one element selected from the group consisting of Mn in a content of 0.05% to 1.0%, Cr in a content of 0.01% to 0.5%, and Zr in a content of 0.01% to 0.2%, with the remainder consisting of Al and inevitable impurities.

For strength and other properties at higher levels, the aluminum alloy may further contain, in mass percent, at least one element selected from the group consisting of Cu in a content of 0.05% to 1.0%, Ti in a content of 0.01% to 0.1%, and Zn in a content of 0.005% to 0.25%. All percentages in contents of individual elements are by mass.

Next, critical significance and preferred ranges of contents of the elements will be illustrated.

Si: 0.7% to 1.5%

Silicon (Si) precipitates, together with Mg, mainly as a needle-like β′ phase in grains by artificial aging and is necessary for imparting higher strength to automobile suspension parts.

Si, if present in an excessively low content, may precipitate in an excessively small amount upon artificial aging and may fail to offer high strength.

In contrast, Si, if present in an excessively high content, may cause coarse elementary Si particles to form and precipitate upon casting and in the course of quenching after solution treatment and may thereby cause the forged aluminum alloy to have lower corrosion resistance and lower toughness. Such a large amount of excessive Si may impede the forged aluminum alloy from having high corrosion resistance, high toughness, and high fatigue properties. In addition, this may also adversely affect hot forgeability and workability to typically cause lower elongation.

On the basis of these, the Si content is controlled in the range of 0.7% to 1.5%.

Mg: 0.6% to 1.2%

Magnesium (Mg) also precipitates, together with Si, as a needle-like β′ phase in grains by artificial aging (aging) and is necessary for imparting higher strength and better ductility to automobile suspension parts.

Mg, if present in an excessively low content, may precipitate in an excessively small amount upon artificial aging and may fail to offer high strength.

In contrast, Mg, if present in an excessively high content, may cause coarse Mg-containing compounds to be formed in grains and at grain boundaries to lower corrosion resistance and toughness. In addition, such excessive Mg may cause the forged aluminum alloy to have excessively high strength (yield strength) at high temperatures, and this may adversely affect hot forgeability and workability.

On the basis of these, the Mg content is controlled in the range of 0.6% to 1.2%.

Fe: 0.01% to 0.5%

Iron (Fe) combines with Si to form intermetallic compounds as dispersed particles (dispersoids), and impedes grain boundary migration ater recrystallization, thereby restrains recrystallization and eliminates or minimizes coarsening of grains. This advantageously contributes to refinement of grains.

In contrast, Fe, if present in an excessively high content, tends to form coarse compounds in grains and at gram boundaries and to cause the forged aluminum alloy to have lower corrosion resistance and toughness. Since such intermetallic compounds formed by Fe readily contain Si, the needle-like β′ phase, which is formed by artificial aging and which requires Si, is decreased. This tends to cause the forged aluminum alloy to have lower strength.

On the basis of these, the Fe content is controlled in the range of 0.01% to 0.5%.

At least one element selected from Mn in a content of 0.05% to 1.0%, Cr in a content of 0.01% to 0.5%, and Zr in a content of 0.01% to 0.2%

As with Fe, manganese (Mn), chromium (Cr), and zirconium (Zr) combine with Si to form intermetallic compounds as dispersed particles (dispersoids), and impede grain boundary migration after recrystallization, thereby restrain recrystallization and eliminate or minimize coarsening of grains. This advantageously contributes to refinement of grains.

In contrast, Mn, Cr, and Zr, if each present in an excessively high content, tend to form coarse compounds in grains and at gram boundaries and to cause the forged aluminum alloy to have lower corrosion resistance and toughness. Since such intermetallic compounds formed by these elements readily contain Si, the needle-like β′ phase, which is formed by artificial aging and which requires Si, is decreased. This tends to cause the forged aluminum alloy to have lower strength.

On the basis of these, the content or contents of at least one of these elements is controlled so that the Mn content falls in the range of 0.05% to 1.0%, the Cr content falls in the range of 0.01% to 0.5%, and the Zr content falls in the range of 0.01% to 0.2%.

At least one element selected from Cu in a content of 0.05% to 1.0%, Ti in a content of 0.01% to 0.1%, and Zn in a content of 0.005% to 0.25%

Copper (Cu), titanium (Ti), and zinc (Z) are equieffective elements to allow the forged material to have strength and toughness at higher levels. When these effects are expected, the forged aluminum alloy may contain one or more of these elements selectively.

Cu offers solid-solution strengthening, thereby contributes to better strength and toughness of the forged material, and effectively significantly promotes age hardening of the final product upon aging. Cu, if present in an excessively low content, may fail to offer these effects on strength improvements. In contrast, Cu, if present in an excessively high content, may cause the forged aluminum alloy microstructure to have significantly high susceptibility (sensitivity) to stress corrosion cracking and to intergranular corrosion and may thereby cause the forged aluminum alloy to have lower corrosion resistance and durability. On the basis of these, the content of Cu, when to be contained, may be controlled in the range of 0.05% to 1.0%.

Zn precipitates and forms Zn—Mg precipitates finely in a high density upon artificial aging and allows the forged aluminum alloy to have better strength and toughness. In addition, solute Zn lowers the potential in grains and causes corrosion not to initiate from grain boundaries, but to be present as general corrosion. This effectively results in reduction of intergranular corrosion and stress corrosion cracking. However, Zn, if present in an excessively high content, may cause the forged aluminum alloy to have remarkably lower corrosion resistance. On the basis of these, the content of Zn, when to be contained, may be controlled in the range of 0.005% to 0.25%.

Ti effectively refines grains of the ingot, allows the forged material microstructure to be fine grains, and allows the forged aluminum alloy to have better strength and toughness. Ti, if present in an excessively low content, may fail to offer these effects. However, Ti, if present in an excessively high content, may form coarse precipitates to lower the workability. On the basis of these, the content of Ti, when to be contained, may be controlled in the range of 0.01% to 0.1%.

It is accepted that the forged aluminum alloy contains other impurity elements as the inevitable impurities constituting part of the remainder of the alloy chemical composition, as long as the impurity elements are in regular amounts according typically to the upper limit specifications in JIS standards, within ranges not adversely affecting properties of the forged aluminum alloy according to the present invention. Such other impurities tend to be contained typically from scrap as a raw material for melting.

For example, impurity elements listed below may be contained up to the after-mentioned contents. Hydrogen is readily contaminated as an impurity and, particularly when the forged material is worked at a low reduction ratio (working ratio), bubbles derived from hydrogen do not undergo compression bonding in working such as forging and cause blisters, which act as fracture origins. This element thereby causes the forged aluminum alloy to have significantly lower toughness and fatigue properties. In particular, the influence of hydrogen is significant typically in suspension parts prepared so as to have higher strength. Accordingly, the hydrogen content is preferably minimized to 0.25 ml or less per 100 g of Al.

Scandium (Sc), vanadium (V), and hafnium (Hf) also tend to be contaminated as impurities and adversely affect the properties of suspension parts. To eliminate or minimize this, the total content of these elements may be controlled to less than 0.3%.

Boron (B), if preset in a content greater than 500 ppm, forms coarse precipitates and thereby lower the workability. On the basis of this, the acceptable content of boron is 500 ppm or less.

Microstructure

On the precondition that the forged aluminum alloy has the above-mentioned alloy chemical composition, the present invention specifies the microstructure of the forged material (forged aluminum alloy) in an observation plane at the center of the thickness (central part of the thickness) of a thickest portion. The forged material is for use to constitute structural components and structural parts of automobiles and other transportation equipment, in particular for use typically in forged automobile suspension parts. The specifying is performed so as to allow the forged material to have strength and ductility both at higher levels (to have both higher strength and better ductility).

Initially, the microstructure is specified so as to have a dislocation density of 1.0×10¹⁴ to 5.0×10¹⁶ per square meter on average, as measured by X-ray diffractometry.

The microstructure is also specified to include small angle grain boundaries with a tilt angle of 2° to 15° in average proportion of 50% or more, where the small angle grain boundaries are present around grains having a misorientation of 2° or more, as measured by SEM-EBSD analysis.

In addition, the microstructure is specified to include precipitates in an average number density of 5.0×10² per cubic micrometer or more, where the precipitates are measurable with a TEM at 300000-fold magnification.

Assume that a forged 6xxx-series aluminum alloy having the alloy chemical composition after solution treatment and quenching is subjected to warm working to be imparted with working strain, and is then subjected to artificial aging. In this case, the resulting forged material (forged aluminum alloy) has strength and ductility both at higher levels, as compared with regular forged materials to which no working strain is imparted.

This is probably because as follows. Heating before the warm working allows uniform, fine β′ phases to precipitate in grains of the forged material. Thereafter the warm working imparts working strain to the forged material to thereby introduce and enhance dislocations. The dislocations restrain β′ phases from precipitating heterogeneously upon artificial aging and thereby allows the forged aluminum alloy to have strength and ductility both at higher levels.

It is also speculated as follows. The uniform, fine β′ phases pin the dislocations introduced by the application of working strain by the action of the subsequent warm working, where the β′ phases have been precipitated in the grains by heating before warm working. This restrains the recovery of dislocations upon artificial aging, ensures work hardening in a sufficient quantity, and contributes to better ductility.

To offer or ensure these advantageous effects, the present invention specifies the average dislocation density, the average proportion of small angle grain boundaries, and the average number density of precipitates, as described above, on the microstructure at the central part of the thickness in a thickest portion of the forged material after artificial aging.

The conditions on the microstructure will be described sequentially below.

Dislocation Density

The present invention specifies and controls the dislocation density in an observation plane at the center of the thickness in a thickest portion of the forged material to be in the range of 1.0×10¹⁴ to 5.0×10¹⁶ per square meter on average, as measured by X-ray diffractometry. This control is performed for higher strength and better ductility of the forged material, in combination with other microstructure controls such as controls on the average proportion of small angle grain boundaries, among grain boundaries, and the average number density of precipitates.

According to the present invention, the forged material after solution treatment and quenching is subjected to warm working to impart working strain (distortion) to the forged material to thereby introduce dislocations again to the forged material. This controls the forged material to have a dislocation density within the specified range. The configuration thus restrains heterogeneous deformation up to a high strain region or up to rupture, where the deformation is caused by the application of external force upon use typically as an automobile suspension part, and allows the forged aluminum alloy to develop excellent work hardening properties (lower yield ratio and higher elongation). This allows the forged aluminum alloy to have high strength in terms of 0.2% yield strength of 400 MPa or more and good ductility in terms of an elongation of 10% or more. This specification (condition) works in combination with other microstructure conditions or controls, such as conditions on the average proportion of small angle grain boundaries in grain boundaries and the average number density of precipitates.

The forged aluminum alloy, if having an excessively low dislocation density less than 1.0×10¹⁴ per square meter, may have inferior work hardening properties as equivalent to conventional forged materials to which no strain is imparted by the warm working. This may cause early rupture in the high strain region upon the application of external force, when the forged aluminum alloy is used typically as an automobile suspension part.

In contrast, the forged aluminum alloy, if having an excessively high dislocation density greater than 5.0×10¹⁶ per square meter, may include smaller amounts of dislocations introduced and accumulated in a high strain region upon the application of external force, when the forged aluminum alloy is used typically as an automobile suspension part. This may also cause early rupture in the high strain region.

Dislocation Density Measurement Method

Although measurement of dislocation density typically with a transmission electron microscope is widely employed, the present invention employs X-ray diffractometry to measure the dislocation density more simply and more reproducibly. Of dislocations, there are “forest” dislocations which are regions (cell walls and shear zones) where linear or streaky dislocations are densely present. The transmission electron microscopic analysis hardly distinguishes such forest dislocations from each other, and this can cause measurement errors in determination of the dislocation density ρ. In contrast, X-ray diffractometry advantageously less causes errors even in analysis of such forest dislocations, because the dislocation density ρ is calculated from half peak widths of diffraction peaks from individual planes in a crystallographic texture, as described below.

When plastic deformation is applied by forging and the warm working to introduce dislocations into a forged material, the resulting forged material has a structure in which lattice distortions are formed as centering around the dislocations. Depending on the arrangement of the dislocations, small angle grain boundaries, cell structures or other structures develop. When such dislocations and domain structures with them are grasped on the basis of X-ray diffraction patterns, distinctive expanses and shapes according to diffraction indices appear in diffraction peaks. The shape (line profile) of such diffraction peaks are analyzed via line profile analysis to determine the dislocation density.

Specifically, the analysis may be performed in the following manner. Initially, the forged material after subjected to solution treatment and quenching is subjected to warm working to be imparted with strain, and is then subjected to artificial aging. From the forged material after artificial aging in a longitudinal section at any position of the thickest portion, three measurement samples (test specimens) including a central part of the thickness are sampled. The samples are, in a word, sliced in parallel with the forged material surface and polished so that the center of the thickness is exposed as an observation plane.

Namely, the term “central part of the thickness” refers to a plane parallel to a forged material surface at the center of the thickness (corresponding to the center of the sheet or plate thickness in a sheet or plate) of the forged material, in a plan view, where the plane extends approximately in parallel with the forged material surface (for example, a horizontal surface) at the center of the thickness.

Of each test specimen, the microstructure of the surface (the plane at the center of the thickness position) is analyzed by X-ray diffractometry, on the basis of which half peak widths of diffraction peaks respectively from (111), (200), (220), (311), (400), (331), (420), and (422) planes (planes of orientations) are determined, where these orientations of planes are principal orientations in the crystallographic texture of the surface portion. The half peak width of the diffraction peak from each plane increases with an increasing dislocation density ρ. Of each test specimen, the rolling surface to be measured by X-ray diffractometry may be as-sampled, or may have undergone cleaning without etching.

Next, a lattice distortion (crystal distortion) ε is determined from the half peak width of the diffraction peak from each plane by Williamson-Hall analysis, on the basis of which the dislocation density ρ can be calculated according to the following expression. The dislocation densities ρ are determined on the three samples sampled from the central parts of the thickness and are averaged to give an average of dislocation density ρ:

ρ=16.1ε² /b ²

where ρ represents the dislocation density; ε represents the lattice distortion; and b represents the magnitude of the Burgers vector. The magnitude of the Burgers vector herein is defined to be 28635×10⁻¹⁰ m.

The Williamson-Hall analysis is a known line profile analysis technique, which is widely used for determining dislocation densities and grain sizes from the relationship between half peak widths and diffraction angles in two or more diffractions. In addition, a series of processes for determining dislocation density by X-ray diffractometry is also publicly known. In the present invention, the dislocation density is referred to as “dislocation density measured by X-ray diffractometry”, where “X-ray diffractometry” herein is employed as a generic name of the series of processes for determining dislocation density by X-ray diffractometry.

Average Proportion of Small Angle Grain Boundaries

To offer higher strength and better ductility of the forged material, the present invention specifies the average proportion of small angle grain boundaries with a tilt angle of 2° to 15° to be 50% or more, where the small angle grain boundaries are present around grains having a misorientation of 2° or more, in an observation plane at the center of the thickness (central part of the thickness) in a thickest portion of the forged material, as measured by electron back-scatter diffraction analysis with canning electron microscopy (SEM-EBSD). The present invention employs this configuration (control) in combination with other microstructure controls such as controls on the average dislocation density and the average number density of precipitates.

The control of the proportion of small angle grain boundaries to be high as in the specified range allows the microstructure to uniformly deform without local concentration or focusing of strain when external force is applied upon use typically as an automobile suspension part. This eliminates or minimizes local rupture and allows the forged material to have high strength in terms of 0.2% yield strength of 400 MPa or more and good ductility in terms of elongation of 10% or more, as employed in combination of other microstructure controls (conditions) such as controls on the average dislocation density and the average number density of precipitates.

In contrast, the forged material, if having an average proportion of small angle grain boundaries less than 50%, does not undergo the mechanism to achieve the high strength and the high elongation, but has a lower elongation, as with conventional forged materials.

As used herein, the term “small angle grain boundary” refers to a grain boundary between grains with a small difference (tilt angle) in crystal orientations of 2° to 15°, out of crystal orientations measured by the after-mentioned SEM-EBSD analysis.

In contrast, the term “large angle grain boundary” refers to a grain boundary between grains having a difference (tilt angle) in the crystal orientations of from greater than 15° to 90°.

In the present invention, the proportion of small angle grain boundaries with a tilt angle of 2° to 15°, which is the average proportion of the small angle grain boundaries, refers to and is defined as the proportion of the total length of the measured small angle grain boundaries (total length of grain boundaries between all measured small angle grains) to the total length of grain boundaries between measured grains with a difference in crystal orientation of 2° to 90° (total length of grain boundaries of all measured grains). Namely, the proportion (%) of the small angle grain boundaries with a tilt angle of 2° to 15° to be specified can be calculated according to the expression: [(Total length of grain boundaries with a tilt angle of 2° to 15°)/Total length of grain boundaries with a tilt angle of 2° to 90°)]×100. The average of the calculated values is controlled to be 50% or more. The upper limit of the proportion of small angle grain boundaries with a tilt angle of 2° to 15° is about 90%, in consideration of limitations in production or hot forging.

Measurement of Average Proportion of Small Angle Grain Boundaries by SEM-EBSD Analysis

The average proportion of small angle grain boundaries with a tilt angle of 2° to 15° around grains with a misorientation of 2° or more in the forged material microstructure is measured by analyzing a microstructure of the forged material after artificial aging by SEM-EBSD analysis, where the microstructure is in an observation plane of the center of the thickness (central part of the thickness) in a thickest portion.

Specifically, the measurement may be performed in the following manner. Three measurement samples including a central part of the thickness are sampled from a longitudinal section at any position of the thickest portion of the forged material after artificial aging, in the same manner as with the sampling of samples for dislocation density measurement. The samples are polished so that an observation plane at the center of the thickness is exposed.

Electron beams are applied at a pitch of 1.0 μm to a measurement region in the observation plane of the forged material using a SEM-EBSD system. The measurement region is a rectangular region having a long side length of 1000 μm and a cross side length of 320 μm.

The average proportion of small angle grain boundaries per each sample is measured in the above manner, and the resulting three measurements of the three samples are averaged (divided by 3).

The SEM-EBSD (EBSP) analysis is a versatile crystal orientation analysis technique using a scanning electron microscope (SEM) including an electron back scattering (scattered) diffraction pattern (EBSD) analysis system.

More specifically, the samples to be observed by SEM-EBSD analysis may be prepared in the following manner. The observation samples (cross-sectional microstructures) are further mechanically polished and then electrically etched to have a mirror surface. Each of the resulting samples is set in a lens barrel of the SEM, and electron beams are applied to the mirror surface of the sample to project an EBSD (EBSP) on the screen. An image of this is taken by a highly sensitive camera and captured as an image into a computer. In the computer, the image is analyzed and compared with patterns obtained by simulation on known crystal systems, and on the basis of the comparison, crystal orientations are determined. The determined crystal orientations are recorded as three-dimensional Eulerian angles typically with position coordinates (x, y, z). This process is automatically performed on all measurement points, and gives crystal orientation data at several tens of thousands to several hundreds of thousands of points upon the completion of measurement. On the basis of the crystal orientation data, grains are distinguished, and the misorientations of grain boundaries are analyzed.

Precipitates

For higher strength and better ductility of the forged material, the present invention specifies the average number density of precipitates to be 5.0×10² per cubic micrometer or more, where the precipitates are measurable with a TEM at 300000-fold magnification in an observation plane at the center of the thickness (central part of the thickness) in a thickest portion of the forged material. This control is employed in combination with the controls on the average dislocation density and the average proportion of small angle grain boundaries.

This allows the forged material to have high strength in terms of 0.2% yield strength of 400 MPa or more and good ductility in terms of elongation of 10% or more, as employed in combination with other microstructure controls (conditions) such as conditions on the average dislocation density and the average proportion of small angle grain boundaries in grain boundaries.

As used herein, the term “precipitates measurable with a TEM at 300000-fold magnification” refers to all precipitates that can be measured identified) with a TEM at 300000-fold magnification, regardless of chemical compositions.

Specifically, the term refers to all precipitates which have granular or massive, rod-like, needle-like, or any other isolated indefinite (complicated) shapes, can be distinguished (identified) typically from the matrix, grain boundaries, and dislocations, and can be observed (determined) by analysis of TEM images.

The minimum size of such precipitates measurable with a TEM at 300000-fold magnification is 5 nm or more in terms of average equivalent circle diameter. Precipitates having a size smaller than this are not measurable and are out of the measurement range.

In this connection, the substantial upper limit in size of the precipitates measurable with a TEM is 1000 nm. This is because automobile suspension part forged materials produced by a common procedure include (are controlled to include) approximately no coarse precipitates having an average equivalent circle diameter greater than 1000 nm, where the coarse precipitates will cause fracture.

As used herein, the “equivalent circle diameter” refers to a circle equivalent diameter obtained by processing images of the identified precipitates, calculating the area of each individual precipitate in the TEM view field, and determining the diameter of a circle having an area equivalent to the area of the precipitate (diameter of the equivalent circle).

The “precipitates” in the present invention are intermetallic compounds having chemical compositions mainly including Mg—Si or Al—Mg—Si—Cu, Al—Mn, Al—Cr, A—Zr, or chemical compositions corresponding to them, except for further including Fe. These intermetallic compounds are derived from the alloy chemical composition and formed upon artificial aging.

The forged material has significantly higher strength (bake hardenability (BH)) when the fine precipitates measurable with the TEM is present (is controlled to be present) in a high average number density of 5.0×10² per cubic micrometer or more.

While its mechanism is still unknown, these precipitates enhance or improve the bake hardenability probably for the following reason. Such transition element-containing dispersed particles having the size and being present in the number density contribute particularly to better work hardening properties upon the application of prestrain, and to restrainment of recovery of dislocations upon artificial aging, where the dislocations are introduced by the application of the prestrain.

In addition, the fine precipitates as above also have an excellent effect of not causing deterioration in elongation of the forged material.

Assume that the forged material includes precipitates measurable with a TEM at 300000-fold magnification at the central part of the thickness of the thickest portion in a low average number density of less than 5.0×10² per cubic micrometer. In this case, the forged material does not undergo the mechanism of bake hardening to achieve the high strength, but has a lower elongation, as with conventional forged materials.

The upper limit of the average number density of precipitates may be about 1.0×10⁵ per cubic micrometer in consideration of limitations in production or hot forging.

Measurement of Average Number Density of Precipitates

The average number density of precipitates as specified in the present invention is measured by measuring or analyzing the microstructure of the forged material after artificial aging in an observation plane at the center of the thickness of the thickest portion with a transmission electron microscope (TEM; such as field-emission transmission electron microscope (FE-TEM)) at 300000-fold magnification.

Specifically, the measurement may be performed in the following manner. Three measurement samples including a central part of the thickness are sampled from the forged material after artificial aging in a longitudinal section at any position in a thickest portion, and thin-film samples for TEM observation are prepared from the measurement samples so that an observation plane at the center of the thickness is exposed.

The TEM thin-film samples are prepared by mechanically polishing the measurement samples so as to have a dimension of 0.05 mm in both thickness directions from the center of the thickness (to have a thickness of 0.1 mm), and thinning the same by twin-et electropolishing into thin films each having a thickness (dimension) of 100 nm from the center of the thickness.

In addition, photos of microstructures of the thin films (samples) are taken using a TEM at 300000-fold magnification, and images thereof are processed, on the basis of which the number of all precipitates which can be identified (distinguished) in the measurement view field is counted, where the total area of the observation view fields is 0.5 μm² or more.

The average number density (number per cubic micrometer) of precipitates in the measurement view field is determined. The average number density measurement is performed on the three samples sampled from the central part of the thickness, the measured three average number densities are averaged, and the average is defined as the average number density (number per cubic micrometer) of precipitates.

As described above, the microstructure and properties of the forged material as specified in the present invention are the microstructure and properties of a forged material obtained by subjecting a forged material after solution treatment and quenching to warm working to impart strain to the forged material, and then subjecting the resulting forged material to artificial aging.

Forged Material Measurement Portion

The measurements of the microstructure and properties are performed at a portion corresponding to the central part of the thickness in a thickest portion of the forged material after artificial aging. When the forged material has a simple shape, so-called “type I”, such as rod-like, plate-like, circular, or cylindrical shape, the central part of the thickness of the forged material to be measured can be specified relative to the center of the forged material.

However, the automobile suspension parts representatively have a complicated shape as follows. This shape is an approximately triangular shape as a whole in a plan view. Ball joints at the three apices of the triangle are coupled to each other though arms. The arms each include ribs and a web, where the ribs are in the periphery and have a narrow width and a large thickness, and the web is in the central portion and has a wide width and a small thickness. The arms each have an approximately H-shaped or approximately U-shaped cross section.

Accordingly, the “central part of the thickness” of such an automobile suspension part is defined herein as the center of the thickness at any position of the thick ribs, where the thick ribs are taken as the thickest portion of the forged material.

Production Method

Next, a method according to the present invention for producing a forged aluminum alloy will be illustrated. The production process itself for the forged aluminum alloy according to the present invention can be performed by a common procedure, in which an aluminum alloy ingot having the chemical composition is subjected sequentially to homogenization and hot forging to give a forged material, and the forged material is subjected sequentially to solution treatment, quenching and artificial aging. Namely, the forged aluminum alloy can be produced without hot extrusion of the ingot.

However, the production method (production process) employs preferred production conditions as described below, such as warm working to be performed after solution treatment and quenching and before artificial aging. Such production conditions are preferred so as to allow the forged aluminum alloy to have the microstructure and to have strength and ductility both at higher levels (higher strength and better ductility) on the precondition that the forged aluminum alloy has high corrosion resistance. The resulting forged aluminum alloy is suitable for use typically as or in an automobile suspension part.

Casting

Casting of a molten aluminum alloy, which is melted and adjusted to have an aluminum alloy chemical composition within the specific range, may be performed by a common melt casting technique. The melt casting technique may be selected as appropriate typically from continuous casting-directed rolling, semicontinuous casting (direct chill (DC) casting), and hot top casting.

However, the casting of the molten aluminum alloy having an aluminum alloy chemical composition within the specific range is preferably performed at an average cooling rate of 100° C./s or more, for refinement of precipitates and decrease of secondary dendrite arm spacing (SDAS).

Homogenization

The homogenization (soaking) of the ingot after casting may be performed by holding the ingot in a temperature range of 450° C. to 580° C. for 2 hours or longer. The homogenization, if performed at a temperature lower than 450° C., may fail to homogenize the ingot due to such excessively low temperature. In contrast, the homogenization, if performed at a temperature higher than 580° C., may cause burning of the ingot surface. Extrusion after homogenization and before hot forging is not necessary, but may be performed when desired.

Hot Forging

After reheating the ingot after homogenization, the hot forging is preferably performed under conditions at a material temperature of from 430° C. to 550° C., a forming die temperature of from 100° C. to 250° C., a minimum reduction of wall thickness of 25% or more, and a maximum reduction of wall thickness of 90% or less.

The hot forging may be performed using a mechanical press or using an oil hydraulic press so as to forge the ingot to a final product shape (or a near net shape) of an automobile suspension part. The hot forging may be performed multiple times as including upset, rough forging, and finish forging, without reheating, or with reheating as needed, during forging.

The hot forging, if performed at a minimum reduction of wall thickness less than 1%, may fail to give the automobile suspension part having the above-mentioned complicated shape with good shape precision by forging, where the minimum reduction of wall thickness is considered as a hot forging reduction ratio. In contrast, the hot forging, if performed at a maximum reduction of wall thickness greater than 90%, may hardly restrain recrystallization and may highly possibly cause coarse recrystallized grains to be formed.

The hot forging, if performed at a forging end temperature after final forging of lower than 300° C., may impede restrainment of recrystallization during forging and solution treatment processes, and this may cause a deformed microstructure to be recrystallized to form coarse grains. These coarse grains, if formed, may impede the forged material from having higher strength and better ductility and may cause the forged material to have lower corrosion resistance, even when the forged material is controlled to have the above-mentioned microstructure. In addition, hot forging, if performed at such a low temperature, may impede refinement of grains in the entire region in a cross section of the forged material. In contrast, the hot forging, if performed at a material temperature higher than 550° C., may highly possibly cause burning of the forged material surface and cause coarse recrystallized grains to be formed.

Solution Treatment and Quenching Treatment

The forged material after the hot forging is subjected to solution treatment and quenching treatment. In the solution treatment, the forged material is preferably held in a temperature range of 530° C. to 570° C. for a time of 1 hour to 8 hours. The solution treatment, if performed at an excessively low temperature and/or for an excessively short time, may become insufficient and may cause insufficient solid-solution of Mg—Si compounds. This may cause the compounds to precipitate in an excessively small amount in the subsequent artificial aging and cause the forged material to have lower strength. The solution treatment may be performed for a long holding time, but may offer saturated effects when performed for a time longer than 8 hours.

After the solution treatment, the forged material may be subjected to quenching preferably at an average cooling rate of 25° C./s or more in a temperature range of from 500° C. down to 100° C. The cooling in the quenching treatment is preferably performed by water cooling, and particularly performed by water cooling (water tank immersion) in which cooling water is circulated with bubbling. This is preferred for ensuring the above-mentioned average cooling rate and for performing homogeneous cooling in which strain on the forged material is eliminated or minimized. The quenching treatment, if performed at an excessively low cooling rate, may cause precipitation typically of Mg—Si compounds and Si at grain boundaries and may thereby cause the product after artificial aging to be susceptible to grain boundary fracture and to have toughness and fatigue properties at lower levels. In addition, such quenching treatment at an excessively low cooling rate may cause Mg—Si compounds and Si, which are stable phases, to be formed also in grains in the course of cooling. This may cause a β phase and a β′ phase to precipitate in smaller amounts upon artificial aging and may cause the forged material to have lower strength.

In contrast, the quenching, if performed at an excessively high cooling rate (if the forged material is cooled excessively rapidly), may cause hardening strain during quenching to be formed in a large amount, and this may disadvantageously require an extra straightening process after quenching, or may cause the straightening process to include steps in a larger number. In addition, such quenching performed at an excessively high cooling rate gives higher (greater) residual stress and may cause the product to have dimensional precision and shape precision at lower levels. In consideration of these, the quenching is preferably performed as hot-water quenching at 30° C. to 85° C. at which temperature quenching strain is relaxed. This is preferred for shortening the product production process and lowering the cost. The hot-water quenching, if performed at a temperature lower than 30° C., may cause greater quenching strain. The hot-water quenching, if performed at a temperature higher than 85° C., may cause the forged material to have toughness, fatigue properties, and strength at lower levels, due to an excessively low cooling rate.

Warm Working

In the present invention, the hot-forged material thus obtained after solution treatment and quenching treatment is subjected to warm working prior to artificial aging so as to be imparted with strain and to be allowed to have the specified microstructure and to have higher strength and better ductility.

The warm working may be performed within 48 hours after solution treatment and quenching treatment. The heating before warm working is preferably performed in a furnace in a temperature range of 140° C. to 220° C. for a placing time (holding time) in the furnace of 20 minutes to 120 minutes. Preferably immediately after the heating, the forged material is subjected to warm working without delay. The holding in the furnace herein is performed by raising the work in temperature for 19 minute to 60 minutes and holding the work at the attained temperature for 1 minute to 60 minutes.

During the heating/holding under the conditions, a uniform fine β′ phase precipitates previously in grains. This heat treatment performed before warm working restrains heterogeneous precipitation of the β′ phase due to dislocations introduced by the subsequent warm working. The already-precipitated β′ phase pins the dislocations, thereby restrains recovery of dislocations upon artificial aging and ensures work hardening in an sufficient magnitude.

In contrast, the heat treatment if performed by placing the work in the furnace for an excessively short time, may cause the β′ phase to little precipitate during temperature rise and heating/holding in the heat treatment before warm working, but may allow dislocations introduced by the subsequent warm working to act as precipitation sites upon artificial aging and to heterogeneously precipitate. In addition, the dislocations may promote diffusion of elements to cause the precipitates to coarsen and to disperse sparsely, and this may possibly cause the forged material to have lower strength.

The warm working is preferably performed at a reduction ratio (working ratio) of 5% to 30%. The warm working, if performed at a reduction ratio less than 5%, may impart a smaller quantity of strain to the forged material, may thereby introduce a smaller amount of dislocations into the forged material, and may fail to offer effective dislocation hardening. In contrast, the warm working, if performed at a reduction ratio greater than 300%, may cause an increased magnitude of accumulated strain to cause the driving force in recovery upon artificial aging to increase. Thus, the warm working less effectively offers strength improvement, because the hardening by dislocation hardening is saturated in amount (magnitude).

In addition, the warm working, if performed to a large magnitude, may increase misorientations of grain boundaries formed by recovery after warm working and/or during artificial aging. This may reduce the proportion of small angle grain boundaries, increase the amount of preferential precipitation at grain boundaries, and cause the forged material to have lower strength contrarily.

The warm working may be performed by a procedure according to the shape of the forged material. When the forged material has a simple shape such as a rod-like, plate-like, circular, or cylindrical shape, the warm working procedure can be selected typically from rolling with rolls and press forming. When the forged material has a complicate shape such as a shape of the automobile suspension part, the warm working procedure can be selected typically from warm closed die forging and warm open die forging.

When a high reduction ratio within the range is desired, the warm working is preferably performed so that the forged material is allowed to have a final product shape by the warm working, namely, the hot forging is performed so that the forged material is allowed to have a near net shape, although this depends on the reduction ratio and working procedure in the warm working.

Artificial Aging

The forged material after the warm working is subjected to artificial aging. To eliminate or minimize the progress of natural aging at room temperature, the artificial aging is preferably performed immediately after the warm working typically within one hour as a rough reference. The artificial aging conditions are preferably selected within a temperature range of 100° C. to 250° C. and within a holding time range of 20 minutes to 8 hours.

However, even within the condition ranges, optimum conditions should be selected corresponding to the chemical composition of the forged material and the conditions of previous processes such as hot forging, solution treatment, quenching treatment, and cold or warm working. Assume that the artificial aging is performed under conditions not corresponding to the chemical composition and the previous process conditions at an excessively low or an excessively high temperature, or for an excessively short holding time. In this case, the artificial aging may impede the forged material from having the desired, specified microstructure and from having high tensile strength, high yield strength, and high elongation.

The homogenization and solution treatment as mentioned above may be performed using an apparatus selected as appropriate typically from air furnaces, induction heating furnaces (induction heaters), and salt-bath furnaces. The artificial aging may be performed using an apparatus selected as appropriate typically from air furnaces, induction heating furnaces, and oil baths.

The forged material according to the present invention, when for use in or as automobile suspension parts, may be subjected to processing such as machining and surface treatment as appropriate before and/or after the artificial aging.

The present invention will be illustrated in further detail with reference to several examples (working examples) below. It should be noted, however, that the examples are by no means intended to limit the scope of the present invention; that various changes and modifications can naturally be made therein without deviating from the spirit and scope of the present invention as described herein; and all such changes and modifications should be considered to be within the scope of the present invention.

EXAMPLES

Next, the present invention will be illustrated with reference to several examples (working examples). Forged materials as materials for automobile suspension parts were produced in the following manner. Initially, hot-forged materials having aluminum alloy chemical compositions given in Table 1 were prepared and subjected to solution treatment and quenching under identical conditions. The hot-forged materials were then subjected sequentially to warm working and artificial aging under individual different conditions given in Table 2 to give the forged materials. The resulting forged materials were subjected to measurements and evaluations on microstructure, mechanical properties, and corrosion resistance as indicated in Table 2.

Specifically, ingots having chemical compositions corresponding to the forged aluminum alloy chemical compositions given in Table 1 were prepared by casting via semicontinuous casting at an average cooling rate of 100° C./s or more, in common in each sample. All the aluminum alloy samples given in Table 1 had a hydrogen content of 0.10 to 0.15 ml per 100 g of Al in common. The symbol “−” in element contents in Table 1 indicates that the content of an element in question is below the detection limit.

In common in each sample, the outer surface of each of the aluminum alloy ingots was faced by a thickness of 3 mm and cut into a round rod-like billet having a length of 120 mm and a diameter of 75 mm. The billet was homogenized at 520° C. for 5 hours and thereafter cooled via forced wind cooling at a cooling rate of 10° C./hr or more using a fan.

The ingot after homogenization was subjected to hot forging, in which forging was performed three times down to a final wall thickness via mechanical press forming using upper and lower dies in common in each sample. The forging was performed under common conditions at a forging start temperature in the range of 500° C. to 520° C., a forming die temperature in the range of 170° C. to 200° C., and a wall thickness change of 75% (greater than 25%) in the central part of the forged material.

In these hot forging operations, each sample was formed into a hot-forged material having a near net shape corresponding to a reduction ratio in after-mentioned warm forging, so as to have a final forged material shape in common in each sample.

These forged materials were, in common in each sample, subjected to solution treatment at 550° C. for 5 hours using an air furnace and then subjected to the water cooling (water tank immersion) at an average cooling rate of 25° C./s or more in the temperature range of from 500° C. down to 100° C.

The hot-forged materials (after solution treatment and quenching treatment) obtained in the above manner were subjected sequentially to warm working and artificial aging under the conditions given in Table 2 to form the individual microstructures.

The warm forging was performed in the following manner. Initially, the forged materials were heated under the heating conditions before warm working given in Table 2 and subjected to warm working at the heating temperature before warm working and the reduction ratio each given in Table 2, via mechanical press forming using upper and lower dies.

The produced forged materials had a suspension part shape in common in each sample. The suspension part shape is an approximately triangular shape as a whole in a plan view. Ball joints at the three apices of the triangle are coupled to each other through arms. The arms each include ribs and a web, where the ribs are in the periphery and have a narrow width and a large thickness (height) of 60 mm, and the web is in the central portion and has a wide width and a small thickness (height) of 31 mm. The arms each have an approximately H-shaped cross section.

The forged materials prepared so as to have the different microstructures as above were subjected to measurements and evaluations on microstructure, mechanical properties, and resistance to intergranular stress corrosion cracking by the following methods. The results of these are presented in Table 2.

Microstructure

The microstructures specified in the present invention were individually measured by the measurement method. Specifically, samples were sampled from a longitudinal section of any central part of the thickness of the rib, which is the thickest portion, of any of the approximately H-shaped arms of the forged material. Of the samples, the average dislocation density (per square meter), the average proportion (%) of small angle grain boundaries with a tilt angle of 2° to 15° around grains having a misorientation of 2° or more, and the average number density of precipitates (number per cubic micrometer) were measured according to the above-mentioned procedures.

Mechanical Properties

A sample was sampled from a central part of the thickness in any portion of the rib, which is the thickest portion, of the forged material. From the sample, three tensile test specimens (L direction) having an outer diameter of 5 mm and a gauge length of 25 mm were prepared so as to include the center of the thickness at a central position of the thickness direction and to have its L direction (longitudinal direction) extending along the longitudinal direction of the forged material. The mechanical properties, such as 0.2% yield strength (MPa) and elongation (%), of the test specimens were measured at room temperature, and the measured values at the three points (three test specimens) were averaged. The tensile speed was set to 5 mm/min at a stress up to the 0.2% yield strength, and to 20 mm/mm at a stress equal to or higher than the 0.2% yield strength.

Acceptance criteria for such forged materials for automobile suspension parts were a 0.2% yield strength of 400 MPa or more and an elongation of 10% or more.

Corrosion Resistance

As the corrosion resistance, resistance to intergranular corrosion (grain boundary corrosion) was evaluated in conformity with the alternate immersion test prescribed in JIS H 8711. Specifically, stress of 300 MPa was applied to a test specimen for stress corrosion cracking resistance evaluation (C-ring test specimen for stress corrosion cracking (SCC) testing), and a time (in day) until intergranular corrosion cracking occurred was measured, regardless of the size of cracking. A sample undergoing intergranular corrosion cracking within a time period shorter than 30 days was evaluated as having poor corrosion resistance (Poor), and a sample undergoing intergranular corrosion cracking within a time period from 30 days to shorter than 60 days was evaluated as having good corrosion resistance (Good).

As clearly demonstrated by Tables 1 and 2, the examples according to the present invention (Examples) had chemical compositions within the ranges specified in the present invention and underwent warm working and artificial aging under conditions within the preferred ranges. As presented in Table 2, the examples have microstructures as specified in the present invention and each have a dislocation density in the range of 1.0×10¹⁴ to 5.0×10¹⁶ per square meter on average as measured by X-ray diffractometry. The examples also have an average proportion of small angle grain boundaries with a tilt angle of 2° to 15° of 50% or more, as measured by SEM-EBSD analysis, where the small angle grain boundaries are present around grains having a misorientation of 2° or more. In addition, the examples have an average number density of precipitates measurable with a TEM at 300000-fold magnification of 5.0×10² per cubic micrometer or more.

As a result, the examples have, as a precondition, excellent corrosion resistance, still have high strength in terms of 02% yield strength of 400 MPa or more and good ductility in terms of elongation of 10% or more, and can combine properties necessary as suspension parts.

In contrast, as in Comparative Examples 13 to 19 in Table 2, samples having alloy chemical compositions within the range corresponding to the alloy number 1 in Table 1, but being produced through warm working under conditions out of the preferred ranges have microstructures at the central part of the thickness, which do not meet conditions specified in the present invention. As a result, these comparative examples have, in common, a 0.2% yield strength and an elongation at significantly lower levels as compared with the examples.

Comparative Example 13 did not undergo warm working before artificial aging.

Comparative Example 14 underwent warm working at an excessively low heating temperature.

Comparative Example 15 underwent warm working at an excessively high heating temperature.

Comparative Example 16 underwent warm working for an excessively short heating-holding time.

Comparative Example 17 underwent warm working for an excessively long heating-holding time.

Comparative Example 18 underwent warm working at an excessively low reduction ratio.

Comparative Example 19 underwent warm working at an excessively high reduction ratio.

Comparative Examples 20 to 23 in Table 2 underwent warm working under conditions within the preferred ranges, but have alloy chemical compositions out of the specified ranges, and have such microstructures at the central part of the thickness as not to meet the conditions on microstructures as specified in the present invention. As a result, these comparative examples have, in common, a 0.2% yield strength and an elongation at significantly lower levels, as compared with the examples.

Comparative Example 20 has a chemical composition corresponding to the alloy number 11 in Table 1 and has a Mg content lower than the lower limit.

Comparative Example 21 has a chemical composition corresponding to the alloy number 12 in Table 1 and has a Si content lower than the lower limit.

Comparative Example 22 has a chemical composition corresponding to the alloy number 13 in Table 1 and does not contain Fe.

Comparative Example 23 has a chemical composition corresponding to the alloy number 14 in Table 1 and contains none of Mn, Cr, and Zr.

These results demonstrate critical significance of the conditions as specified in the present invention on chemical compositions and microstructures to give forged 6xxx-series aluminum alloys having excellent corrosion resistance and still having both high strength and good ductility.

TABLE 1 Chemical composition of forged 6xxx series aluminum alloy Alloy (in mass percent, the remainder including Al) number Mg Si Fe Mn Cr Zr Cu Ti Zn 1 0.8 1.0 0.25 — 0.08 — — 0.02 — 2 0.8 1.0 0.14 0.20 — 0.03 — — — 3 0.7 1.1 0.14 0.25 0.03 0.02 0.4 0.02 — 4 1.2 1.1 0.14 0.08 0.15 — 0.3 0.02 0.05 5 1.1 0.8 0.14 0.46 — — — 0.02 — 6 0.7 1.3 0.02 0.25 0.02 — 0.3 0.04 — 7 0.6 1.2 0.14 0.06 0.40 — — 0.02 — 8 0.7 1.1 0.14 0.15 0.03 0.18 0.4 0.01 — 9 0.7 1.1 0.14 0.25 0.03 — 0.7 0.02 — 10 0.7 1.1 0.14 0.25 0.03 — 0.4 0.02 0.20 11 0.4 1.0 0.18 0.08 — — 0.1 0.02 — 12 1.0 0.5 0.20 0.15 0.03 0.02 0.1  0.015 0.1  13 0.6 0.7 — 0.08 — — 0.1 0.02 — 14 0.6 0.9 0.25 — — — 0.1  0.025 —

TABLE 2 Production method of forged 6xxx series aluminum alloy Warm working conditions after solution treatment Heating conditions Artificial aging conditions Alloy before warm working after warm working number in Temperature Placing time Warm working ratio Temperature (° C.) Category No. Table 1 (° C.) (min) (%) for time (hr) Examples 1 1 160 100 30 170° C. for 8 hr 2 2 180 60 20 190° C. for 1 hr 3 3 140 110 15 180° C. for 3 hr 4 4 220 40 20 230° C. for 0.4 hr 5 5 190 20 25 200° C. for 1 hr 6 6 170 120 10 180° C. for 2 hr 7 7 180 90 5 190° C. for 2 hr 8 8 170 80 30 200° C. for 0.5 hr 9 9 200 30 15 220° C. for 0.5 hr 10 10 150 120 10 180° C. for 5 hr Comparative 13 1 — — — 190° C. for 4 hr Examples 14 1 120 30 5 190° C. for 4 hr 15 1 250 90 10 190° C. for 4 hr 16 1 140 5 5 160° C. for 4 hr 17 1 220 180 10 190° C. for 4 hr 18 1 150 20 2 190° C. for 4 hr 19 1 140 30 50 190° C. for 4 hr 20 11 160 60 10 190° C. for 4 hr 21 12 160 60 10 190° C. for 4 hr 22 13 160 60 10 190° C. for 4 hr 23 14 160 60 10 190° C. for 4 hr Properties of forged 6xxx series aluminum alloy after artificial aging Average proportion Average Resistance Dislocation of small number 0.2% to Alloy density angle grain density of Yield Elonga- intergranular number in (average) boundaries precipitates stength tion corrosion Category No. Table 1 ×10¹⁵/m² % ×10²/μm³ MPa % cracking Examples 1 1 4.71 78 35.1 429 16 Good 2 2 2.45 81 41.3 414 17 Good 3 3 2.13 72 29.6 411 17 Good 4 4 1.32 69 18.1 403 16 Good 5 5 1.18 64 12.8 401 17 Good 6 6 3.42 75 33.4 425 16 Good 7 7 1.39 58 26.3 408 17 Good 8 8 1.85 74 24.9 410 17 Good 9 9 1.13 63 46.6 416 16 Good 10 10 2.98 71 31.9 421 16 Good Comparative 13 1 0.03 47 4.6 364 18 Good Examples 14 1 0.42 63 3.6 383 18 Good 15 1 0.53 54 2.9 378 17 Good 16 1 0.07 68 1.6 366 9 Good 17 1 0.65 61 3.7 372 17 Good 18 1 0.05 49 4.6 364 18 Good 19 1 1.42 64 4.7 391 17 Good 20 11 0.31 60 3.7 377 16 Good 21 12 0.18 67 2.5 352 17 Good 22 13 0.09 52 5.4 386 16 Good 23 14 0.06 53 5.2 390 16 Good

The present invention can provide forged 6xxx-series aluminum alloys having excellent corrosion resistance and still having both high strength and good ductility. The present invention therefore enlarges applications of hot-forged 6xxx-series aluminum alloys to transportation equipment such as automobile suspension parts and has significant industrial value. 

1: A forged aluminum alloy having excellent strength and ductility, the forged aluminum alloy comprising, in mass percent: Si in a content of 0.7% to 1.5%; Mg in a content of 0.6% to 1.2%; Fe in a content of 0.01% to 0.5%; at least one element selected from the group consisting of: Mn in a content of 0.05% to 1.0%; Cr in a content of 0.01% to 0.5%; and Zr in a content of 0.01% to 0.2%; and Al and inevitable impurities, wherein: the forged aluminum alloy having a microstructure in an observation plane at a center of a thickness in a thickest portion of the forged aluminum alloy; the microstructure has a dislocation density of from 1.0×10¹⁴ to 5.0×10¹⁶ per square meter on average as measured by X-ray diffractometry; the microstructure comprises small angle grain boundaries with a tilt angle of 2° to 15° in an average proportion of 50% or more as measured by SEM-EBSD, the small angle grain boundaries being present around grains having a misorientation of 2° or more; and the microstructure comprises precipitates in an average number density of 5.0×10² per cubic micrometer or more, the precipitates being measurable by a transmission electron microscope (TEM) at 300000-fold magnification. 2: The forged aluminum alloy according to claim 1, further comprising, in mass percent, at least one element selected from the group consisting of: Cu in a content of 0.05% to 1.0%; Ti in a content of 0.01% to 0.1%; and Zn in a content of 0.005% to 0.25%. 3: The forged aluminum alloy according to claim 1, which has a tensile strength of 420 MPa or more, a 0.2% yield strength of 400 MPa or more, and an elongation of 10% or more. 4: A method for producing a forged aluminum alloy having excellent strength and ductility, the method comprising: preparing an aluminum alloy ingot comprising, in mass percent: Si in a content of 0.7% to 1.5%; Mg in a content of 0.6% to 1.2%; Fe in a content of 0.01% to 0.5%; and at least one element selected from the group consisting of: Mn in a content of 0.05% to 1.0%; Cr in a content of 0.01% to 0.5%; and Zr in a content of 0.01% to 0.2%, and Al and inevitable impurities; subjecting the aluminum alloy ingot sequentially to homogenization and hot-forging to give a forged material; and subjecting the forged material sequentially to solution treatment, quenching, warm working, and artificial aging in the specified sequence, wherein: the forged material after artificial aging has a microstructure in an observation plane at a center of a thickness in a thickest portion of the forged material; the microstructure has a dislocation density of from 1.0×10¹⁴ to 5.0×10¹⁶ per square meter on average as measured by X-ray diffractometry; the microstructure comprises small angle grain boundaries with a tilt angle of 2° to 15° in an average proportion of 50% or more as measured by SEM-EBSD, the small angle grain boundaries being present around grains having a misorientation of 20 or more; and the microstructure comprises precipitates in an average number density of 5.0×10² per cubic micrometer or more, the precipitates being measurable by a TEM at 300000-fold magnification. 